Advanced Fe-5Cr-X Alloy

ABSTRACT

A tubular article can be formed from high temperature steam oxidation resistant and high temperature creep resistant alloy steel. The steel can include a chemical composition that include Fe, C, Si, Mn, Ni, Cr, Cu, Ti, Nb, Mo, and N, and optionally other elements. The steel alloy can include 0.06 to 0.15 wt % C, 0.1 to 0.5 wt % Si, 0.2 to 0.6 wt %, 0.05 to 0.4 wt % Ni, 4.5 to 6.0 wt % Cr, 1.0 to 2.0 wt % Cu, 0.04 to 0.08 wt % Ti, 0.01 to 0.06 wt % Nb, 0.45 to 1.2 wt % Mo, and 0.008 to 0.05 wt % N, up to 0.01 wt % of optional element Al, up to 0.01 wt % of optional element Zr, up to 3.0 wt % of optional element Co, up to 0.07 wt % of optional element V, up to 3.0 wt % of optional element W, up to 0.015 wt % of optional element P, up to 0.003 wt % of optional element S, up to 0.1 wt % of optional element Ca, up to 0.1 wt % of optional element Ta, up to 0.1 wt % of optional element Mg, up to 0.1 wt % of optional element Se, up to 0.1 wt % of optional element Te, up to 0.1 wt % of optional element B, up to 0.1 wt % of optional element Bi, and Fe. The steel can include copper precipitates and fine carbides, nitrides, or both. The steel can have a final microstructure comprising tempered martensite, tempered bainite, or a combination thereof. The steel can contain less than 2 vol % residual austenite.

STATEMENT OF GOVERNMENT INTEREST

This invention was made with government support under Agreement No. NFE-07-01188 and U.S. Department of Energy Prime Contract No. DE-ACO5-00OR22725. The government has certain rights in the invention.

TECHNICAL FIELD

This invention relates to a tubular article formed from an alloy steel with high temperature steam oxidation and high temperature creep resistance, as well as a method of producing the tubular article. More particularly, the tubular articles formed from the steel are suitable for boilers, pipes, heat exchangers, steam generators, water panels, waterwall tubes, large heaters/collectors, and the like for power-generating equipment and plants.

BACKGROUND

The efficiency of thermal power plants is currently limited by the long-term creep strength and the steam oxidation resistance of the commercially available ferritic/martensitic steel grades. Low alloy ferritic steel grades (e.g., containing 1 wt % Cr) are typically used for the construction of boiler waterwalls. However, a new generation of power plants will be operating at even higher pressures and temperatures under atmospheres containing steam to increase efficiency. Accordingly, important requirements on the materials, from both the mechanical and chemical stability perspective need to be imposed. As such, alloy steels with a combination of characteristics related to yield strength, high resistance to corrosion and steam oxidation, high creep strength, weldability, ability to withstand hot- and cold-bending and expansion, ability to be continuously cast into billets (i.e. hot ductility) are desirable.

Extensive research has been done to improve the steam oxidation resistance (for oxidation rates up to 600° C.) by increasing the Cr content. (Materials Chemistry and Physics 141 (2013) p. 432-439, Alina Agiiero, Vanessa González, Peter Mayr, Krystina Spiradek-Hahn). Hereby, it has been shown that Cr wt % higher than 9 is required; on the other hand such high Cr content is detrimental to the creep strength. FIG. 1 shows the improvement in steam oxidation resistance of 9-12 wt % Cr materials over low alloy ferritic steel grades.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows bore oxide scale thickness as a function of exposure time and temperature for ferritic (1-2 wt % Cr) and martensitic (9-12 wt % Cr) steels (after Sumitomo Metal Industries Ltd, “Steam Oxidation on Cr—Mo-Steel Tubes”, Paper No. 805, 1443A, 1989).

FIG. 2A shows variation of Vickers hardness with tempering temperature for alloy steels 1-3. FIG. 2B shows variation of Vickers hardness with tempering temperature for alloy steels 4-6.

FIGS. 3A-3F show optical micrographs of alloy steels 1-6, respectively, after normalizing at 1050° C. and tempering at 700° C.

FIG. 4A shows variations of temperature for alloy steel 1 and 3 specimens with time during Gleeble testing. FIG. 4B shows variations of axial displacement with temperature of alloy steel 1 and 3 specimens during Gleeble testing.

FIG. 5A shows continuous cooling transformations (CCTs) of alloy steel 3 and 7 specimens. FIGS. 5B-5D show optical micrographs of alloy steel 7 specimens cooled at 1° C./s, 60° C./s, and 0.5° C./s, respectively.

FIG. 6 shows phase amount as a function of cooling rate for an alloy steel 7 specimen.

FIG. 7 shows a secondary phase amount calculation for alloy steel 7.

FIG. 8A shows a thin film transmission electron microscopy (TEM) image of incoherent Cu precipitates in an alloy steel 7 specimen. FIGS. 8B and 8C show energy-dispersive X-ray (EDS) spectra for an alloy steel 7 specimen.

FIG. 9A shows an extraction replica TEM image of carbide precipitates in an alloy steel 7 specimen. FIG. 9B shows a M₂₃C₆ selective area diffraction pattern for an alloy steel 7 specimen.

FIGS. 10A and 10B show X-ray diffraction (XRD) patterns from normalized and tempered specimens of alloy steel 7.

FIG. 11 shows the chromium content in solid solution in alpha ferrite matrix at various tempering temperatures, calculated under equilibrium conditions for alloy steel 7.

DETAILED DESCRIPTION

Tubular articles formed of high alloy steel described herein have high temperature steam oxidation and high temperature creep resistance.

As described herein, “tubular article” generally refers to a tube or pipe and includes seamless tubular articles. The tubular articles described herein have various dimensions. In some examples, the tubular articles have a wall thickness (WT) between 2 mm and 30 mm. In other examples, the tubular articles have an outer diameter (OD) between 20 mm and 624 mm. These steam oxidation and creep resistant tubular articles described herein are suitable for use in environments with high service temperature, up to 625° C., including boilers, pipes, heat exchangers, steam generators, water panels, waterwall tubes, large heaters/collectors, and the like for power-generating equipment and plants.

The steel used to form the tubular articles includes:

-   -   C of 0.06 to 0.15 wt %;     -   Si of 0.1 to 0.5 wt %;     -   Mn of 0.2 to 0.6 wt %;     -   Ni of 0.05 to 0.4 wt %;     -   Cr of 4.0 to 6.0 wt %,     -   Cu of 1.0 to 2.0 wt %;     -   Ti of 0.04 to 0.08 wt %;     -   Nb of 0.01 to 0.06 wt %;     -   Mo of 0.45 to 1.2 wt %;     -   N of 0.008 to 0.05 wt %;     -   Al of less than or equal to 0.01 wt %;     -   Zr of less than or equal to 0.01 wt %;     -   Co of less than or equal to 3.0 wt %;     -   V of less than or equal to 0.07 wt %;     -   W of less than or equal to 3.0 wt %;     -   P of less than or equal to 0.015 wt %;     -   S of less than or equal to 0.005 wt %;     -   and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or         equal to 0.1 wt %;     -   and a balance comprising Fe and impurities.

In one example, wt % Ta+wt % Nb<0.1 wt %.

In another example, less than 4 wt % of Cr is kept in solid solution at room temperature.

In some cases, the steel used to form the tubular articles consists essentially of:

-   -   C of 0.06 to 0.15 wt %;     -   Si of 0.10 to 0.5 wt %;     -   Mn of 0.2 to 0.6 wt %;     -   Ni of 0.05 to 0.25 wt %;     -   Cr 5.0 to 6.0 wt %;     -   Cu of 1.0 to 2.0 wt %;     -   Ti of 0.01 to 0.10 wt %;     -   Nb of 0.02 to 0.10 wt %;     -   Mo of 0.45 to 1.2 wt %;     -   N of 0.008 to 0.05 wt %;     -   Al of less than or equal to 0.01 wt %;     -   Zr of less than or equal to 0.01 wt %;     -   Co of less than or equal to 3.0 wt %;     -   V of less than or equal to 0.07 wt %;     -   W of less than or equal to 3.0 wt %;     -   P of less than or equal to 0.015 wt %;     -   S of less than or equal to 0.005 wt %;     -   and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or         equal to 0.1 wt %;     -   and a balance comprising Fe and impurities.

In some cases, the steel used to form the tubular articles consists essentially of:

-   -   C of 0.06 to 0.15 wt %;     -   Si of 0.10 to 0.5 wt %;     -   Mn of 0.2 to 0.6 wt %;     -   Ni of 0.05 to 0.25 wt %;     -   Cr 5.0 to 6.0 wt %;     -   Cu of 1.0 to 2.0 wt %;     -   Ti of 0.01 to 0.10 wt %;     -   Nb of 0.02 to 0.10 wt %;     -   Mo of 0.45 to 1.2 wt %;     -   N of 0.008 to 0.05 wt %;     -   Al of less than or equal to 0.01 wt %;     -   Zr of less than or equal to 0.01 wt %;     -   Co of less than or equal to 3.0 wt %;     -   V of less than or equal to 0.07 wt %;     -   W of less than or equal to 3.0 wt %;     -   P of less than or equal to 0.015 wt %;     -   S of less than or equal to 0.005 wt %;     -   and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or         equal to 0.1 wt %;     -   and a balance comprising Fe and impurities.

In certain cases, the steel used to form the tubular articles consists essentially of:

-   -   C of 0.06 to 0.15 wt %;     -   Si of 0.10 to 0.5 wt %;     -   Mn of 0.2 to 0.6 wt %;     -   Ni of 0.05 to 0.25 wt %;     -   Cr of 4.0 to 6.0 wt %;     -   Cu of 1.0 to 2.5 wt %;     -   Ti of 0.04 to 0.08 wt %;     -   Nb of 0.02 to 0.10 wt %;     -   Mo of 0.45 to 1.2 wt %;     -   N of 0.008 to 0.05 wt %;     -   Al of less than or equal to 0.01 wt %;     -   Zr of less than or equal to 0.01 wt %;     -   Co of less than or equal to 3.0 wt %;     -   V of less than or equal to 0.07 wt %;     -   W of less than or equal to 3.0 wt %;     -   P of less than or equal to 0.015 wt %;     -   S of less than or equal to 0.005 wt %;     -   and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or         equal to 0.1 wt %;     -   and a balance comprising Fe and impurities.

In certain cases, the steel used to form the tubular articles consists of:

-   -   C of 0.06 to 0.15 wt %;     -   Si of 0.10 to 0.5 wt %;     -   Mn of 0.2 to 0.6 wt %;     -   Ni of 0.05 to 0.25 wt %;     -   Cr of 4.0 to 6.0 wt %;     -   Cu of 1.0 to 2.5 wt %;     -   Ti of 0.04 to 0.08 wt %;     -   Nb of 0.02 to 0.10 wt %;     -   Mo of 0.45 to 1.2 wt %;     -   N of 0.008 to 0.05 wt %;     -   Al of less than or equal to 0.01 wt %;     -   Zr of less than or equal to 0.01 wt %;     -   Co of less than or equal to 3.0 wt %;     -   V of less than or equal to 0.07 wt %;     -   W of less than or equal to 3.0 wt %;     -   P of less than or equal to 0.015 wt %;     -   S of less than or equal to 0.005 wt %;     -   and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or         equal to 0.1 wt %;     -   and a balance comprising Fe and impurities.

The steel includes fine carbides. The fine carbides may include, for example, one or more of fine titanium, niobium, molybdenum and vanadium carbides. As described herein, “fine carbides” generally refer to carbides having a mean diameter of less than 200 nm. In some examples, the fine carbides have a mean diameter between 5 nm and 100 nm. In some cases, the fine carbides have a size distribution with a mean diameter of 65 nm or less.

The steel also includes copper precipitates. The copper precipitates typically have a mean diameter of less than 40 nm.

The steel may also include fine nitrides. As described herein, “fine nitrides” generally refer to nitrides having a mean diameter of less than 200 nm. In some examples, the fine nitrides have a mean diameter between 5 nm and 100 nm.

The steel has a final microstructure including tempered martensite, tempered bainite, or a combination thereof. The steel may include at least 60 vol %, at least 70 vol %, at least 80 vol %, at least 90 vol %, or at least 95 vol % of tempered martensite. The steel may include up to 5 vol %, up to 10 vol %, up to 15 vol %, up to 20 vol %, or up to 25 vol % of tempered bainite. The steel is substantially free of austenite. As used herein, “substantially free of austenite” generally refers to less than 2 vol % austenite. In some cases, the steel includes less than 1 vol %, less than 0.5 vol %, or less than 0.1 vol % austenite. In certain cases, the steel includes ferrite of less than 10 vol % or less than 5 vol %. In certain cases, the steel consists essentially of tempered martensite, tempered bainite, or a combination thereof.

The steel described herein has a yield strength exceeding 450 MPa (e.g., yield strength between 570 MPa and 640 MPa), a tensile strength exceeding 585 MPa at room temperature, a hardness of 200 to 265 Vickers (HV), and a creep strength of about 70 MPa for 100,000-hour rupture at 600° C. The steel has a minimum absorbed energy in Charpy V-notch impact test of 150 J at 0° C. The steam oxidation resistance of the alloy steel is better than that of ASTM A213 T22 and ASTM A213 T23 grade steel by 20-30%. The alloy steel has good weldability with no post-weld heat treatment (PWHT) of butt welds and submerged arc welding (SAW) of tube to strip (300 HV maximum hardness). The alloy steel is also able to withstand hot- and cold-bending/expansion, be continuously cast into round billets (e.g. 148 mm diameter), and has good high-temperature ductility. The heat treatment includes normalization and tempering.

In the following description, the purpose of each of the chemical elements contained in the alloy steel is discussed, and the reason why the values for each element should be restricted to a specific range is also discussed. The content of each chemical element is represented in terms of the percentage by weight (wt %).

Carbon (C) combines with Cr, Mo, V, Ti and Nb to form carbides, which result in improved high temperature creep strength and increased microstructural stability under prolonged exposures at elevated temperature. Also, with reduced carbon content, the ferrite formation during air-cooling is promoted, degrading the strength due to the decreased amount of martensite upon normalizing treatment. With increased carbon content, the A_(C1) transformation point decreases, reducing applicability for high temperature service. Moreover, if C content exceeds 0.15%, carbides coarsen more quickly, resulting in the reduction in the creep strength at high temperature. Additionally, an increase in the amount of C may lead to undesirably high level of hardness and poor weldability. Therefore, the carbon content should be limited within the range 0.06-0.15 wt %, preferably 0.07-0.14 wt %, and most preferably 0.07-0.12 wt %.

Manganese (Mn) is an element useful for fixing S in molten steel (by forming MnS) to give improved hot working property. Mn content can be adjusted to improve steel hardenability and strength. However, Mn contents greater than 0.6% promote segregation and therefore uneven strength, decrease weldability and result in low A_(C1) transformation point and decreased creep resistance. Therefore, the Mn content should be limited within the range 0.2-0.60 wt %, and preferably 0.3-0.50 wt %.

Silicon (Si) is added as a deoxidizing agent and to increase resistance to steam oxidation resistance. With a silicon content under 0.1 wt %, the desired effect cannot be obtained. However, with a silicon content over 0.6 wt %, the amount of ferrite in the steel increases, thus leading to lower (creep) strength. Therefore, the silicon content should be limited within the range 0.10-0.60 wt %, and preferably 0.15-0.40 wt %, and most preferably 0.20-0.35 wt %.

Chromium (Cr) is added to give the steel a satisfactory level of hot corrosion resistance, and is an essential element in the formation of a stable oxide scale for resistance to high-temperature oxidation, especially steam oxidation. Cr also forms carbides which improve creep strength of steel by means of precipitation hardening. In order to obtain these effects, the amount of chromium in steel is 4.0 wt % or more. When the amount is over 6.0 wt %, weldability deteriorates.

Therefore, Cr content should be limited within the range 4.0-6.0 wt %, and preferably 4.5-6.0 wt %, and most preferably 4.5-5.3 wt %.

Molybdenum (Mo) is added to achieve solution strengthening and improved creep strength because it is effective in making carbides that remain stable at high temperature for a long period of time. With a molybdenum content under 0.45 wt %, the desired effect cannot be obtained. However, with a molybdenum content over 1.2 wt %, precipitation of Laves phase (Fe₂Mo), that significantly reduces toughness and creep strength in the long term, can occur during service at high temperature. Additionally, increases in molybdenum content have a significant impact on cost. Therefore, the molybdenum content should be limited within the range 0.45-1.2 wt %, preferably 0.45-1.0 wt %, most preferably 0.45-0.70 wt %.

Titanium (Ti) produces stable carbides and nitrides that result in higher creep strength and long term microstructural stability. However, with a titanium content under 0.01 wt % the desired effect cannot generally be obtained, and with a titanium content over 0.10 wt %, very coarse carbides may result in the steel matrix. The formation of these large carbides is deleterious to mechanical properties as they act as stress concentration risers and reduce creep strength as the number of fine carbo-nitrides is reduced. Therefore, Ti content should be limited to within the range 0.01-0.10 wt %, and preferably 0.02-0.08 wt %, and most preferably 0.04-0.06 wt %.

Niobium (Nb) forms fine carbides and carbo-nitrides that increase the high temperature creep strength. This element also promotes grain growth control during normalization treatment. Thus, at least 0.02 wt % of Nb is required. Over 0.10 wt %, however, Nb may cause the precipitation of coarse carbides and carbo-nitrides reducing ductility and creep strength. Moreover, Nb impairs weldability and more severe post-welding heat treatments are required to relieve internal stresses and reduce hardness. For this reason, Nb content should be limited within the range 0.02-0.10 wt %, preferably 0.02-0.08 wt %, most preferably 0.03-0.06 wt %.

Vanadium (V) combines with carbon and nitrogen to form finely dispersed precipitates such as V(C,N), which are stable at relatively high temperature for an extended period of time. The dispersed V(C, N) is effective for improving long-term creep strength. With a vanadium content below 0.005 wt %, the desired effect cannot be obtained. With a vanadium content over 0.20 wt %, toughness and creep strength can be impaired. Also, similarly to Nb, an excessive amount of V impairs weldability because more severe post-welding heat treatments are required to relieve internal stresses and reduce hardness. Therefore, V content should be limited within the range 0.005-0.20 wt %, and preferably 0.005-0.12 wt %, and most preferably 0.005-0.07 wt %.

Nitrogen (N) combines with titanium, niobium and vanadium to form fine nitrides, which are effective to strengthen the steel and improve its creep resistance. These compounds are commonly described as MX-phases, where M denotes a metal atom and X denotes nitrogen or carbon. In fact, as carbides can be also present as MC and carbides and nitrides have mutual solubility, complex carbo-nitrides denoted as (Nb,Ti)(C,N) can form. Nitrogen added to steels can promote strengthening beyond that is gained by carbon alone. In addition, the particles containing more nitrogen and higher Ti and Nb amounts, are more stable than the others during long exposure at high temperature (ageing), thereby maintaining the strengthening effect over time (e.g., when the steel is used for elevated-temperature service). If nitrogen content is below 0.008 wt % an important strengthening effect cannot be obtained. With a nitrogen content over 0.035 wt %, coarse nitrides less effective in precipitation strengthening form and weldability can be decreased. Therefore, nitrogen content should be limited to less than 0.035 wt %, and preferably less than 0.02 wt %.

Nickel (Ni) is an austenite stabilizer and may be added to effectively obtain a martensitic structure after quenching. However, at relatively low cooling rates it promotes formation of retained austenite, which is detrimental to creep strength. Moreover, the A_(C1) transformation point is significantly lowered and creep strength may be reduced increasing Ni content. Additionally, increasing in nickel content has a significant impact on costs. Therefore, the Ni content should be limited to about 0.25 wt %, preferably 0.15 wt % or less. Moreover, current steelmaking processes and procedures do not allow to maintain the Ni content consistently below 0.05% at limited cost. Therefore the minimum Ni content is 0.05 wt %.

Phosphorus (P) and Sulfur (S) are unavoidable impurities adversely affecting steel properties such as toughness and weldability. Phosphorus should be limited to an amount less than 0.025 wt %, preferably less than 0.015 wt %, most preferably less than 0.012 wt %. Sulfur should be limited to an amount less than 0.025 wt %, preferably less than 0.015 wt %, most preferably less than 0.003 wt %.

Aluminum (Al) is usually added as a deoxidizing agent to liquid steel. However, when the aluminum content is over 0.02 wt %, toughness and high-temperature creep strength can be deteriorated. Therefore, the aluminum content should be limited to about 0.02 wt % or less, and preferably 0.01 wt % or less. Because it is very difficult and expensive to reduce the Al content at extremely low levels, the minimum Al content in deoxidized (killed) steel is set at 0.005%.

Copper (Cu), like Mn and Ni, is an element that stabilizes the austenite, and contributes to strengthening through precipitation hardening by forming fine ε-Cu particles during tempering. Below 1.0 wt %, this element does not provide this effect because it remains in solid solution, and above 2.5%, it may reduce the hot workability due to liquation at grain boundaries. Thus, the preferred Cu content should be at least 1.0% and at most 2.5%. More preferably, it should be at least 1.5 wt % and at most 2.0 wt %.

Tungsten (W) is an element able to increase the creep rupture strength of this steel by precipitation of carbides and solution hardening. This element may form a solid solution in the alloy steel, with consequent reinforcement of the matrix. It promotes also formation of fine and dispersed phases, increasing the high temperature strength and the high temperature creep strength. W content greater than 0.005 wt % is required to appreciate strength increase. However, over 2.0 wt % W may produce coarse Laves phases (Fe₂W) even after relative short time at high temperature. For this reason, W content should be at least 0.005 wt % and at most 3.0 wt %. The upper limit of 2.0 wt % is preferred. More preferably, the upper limit should be 1.5 wt %. When W is added in combination with Mo, the above concentration limits apply to (Mo+0.5 W) content, i.e. Mo+0.5 W should be at least 0.005 wt:% and at most 3.0 wt %. The upper limit of 2.0 wt % is preferred. More preferably, the upper limit of Mo+0.5 W should be 1.5 wt %.

Cobalt (Co) slows diffusion processes that cause coarsening of secondary phase particles, and thus decreases loss of strength after long term exposure at high temperature. In order to exploit this effect, at least 0.05% of Co is required. Over 3.0%, Co may reduce the ductility and as it is an expensive alloying element its maximum content should be limited to at most 3.0 wt %, preferably 2 wt %, and most preferably 1.5 wt %.

Tantalum (Ta) plays the same role as Nb, forming fine carbides and carbo-nitrides that increase the high temperature creep strength. Ta may replace Nb with similar effects.

In order to be effective, at least 0.02 wt % of Ta is required. Over 0.10 wt %, however, Ta may cause the precipitation of coarse carbides and carbo-nitrides reducing ductility and creep strength, and negatively affecting weldability. For this reason, Ta content should be limited within the range 0.02-0.10 wt %, preferably 0.02-0.08 wt %, most preferably 0.03-0.06 wt %. When Ta is added in combination with Nb, the above concentration limits apply to (Ta+Nb) content, i.e. Ta+Nb should be limited within the range 0.02-0.10 wt %, preferably 0.02-0.08 wt %, most preferably 0.03-0.06 wt %.

Boron (B) works as an effective inhibitor that prevents the growth of precipitated carbides, carbo-nitrides and Laves phases. When combined with elements such as Nb, Mo and others, B may increase the high temperature creep strength. Also, B improves hardenability, and thus allows obtaining a ferrite-free microstructure in thick wall products, even with still air cooling. Below 0.0005 wt %, this element does not provide any significant effect. However, B may combine with nitrogen, forming BN nitride that may reduce the high temperature creep ductility and toughness. For this reason, maximum B content should be limited to 0.01 wt %, preferably 0.005 wt %.

Calcium (Ca) is an element whose addition to the steel composition may assist with control of the shape of inclusions by forming fine and substantially round sulfides. In an embodiment, in order to provide these benefits, the Ca content of the steel composition is selected to be greater than or equal to about 0.0010 wt % when the sulfur content of the steel composition is higher than about 0.0020 wt %. However in other embodiments, if the Ca content of the steel composition exceeds about 0.0060 wt %, the effect of the Ca addition may be saturated and the risk of forming clusters of Ca-rich non-metallic inclusions that reduce ductility and creep resistance may be increased. Accordingly, in certain embodiments, the maximum Ca content of the steel composition is selected to be less than or equal to about 0.0060 wt %, preferably less than or equal to about 0.0050 wt %, and most preferably less than or equal to about 0.0040 wt %, while the minimum Ca content is selected to be greater than or equal to about 0.0010 wt %.

Magnesium (Mg) may be present as an unavoidable residual element. It can be found together with Ca and Si, forming mixed oxides that remain in the form of non-metallic inclusions. While this element cannot be systematically reduced to below 0.0010 wt % with current steelmaking technologies with affordable costs, it should be maintained below the 0.0050 wt % threshold to guarantee a good impact toughness level.

Arsenic (As), Antimony (Sb), Tin (Sn), Selenium (Se), Tellurium (Te), Lead (Pb), Bismuth (Bi), and Zirconium (Zr), are all elements that are not intentionally added. However, depending on the manufacturing process, the presence of these impurity elements may be unavoidable.

Some of these impurities need to be limited to very low amount in order to inhibit embrittlement phenomena. Therefore, the maximum allowable content of these single elements should be 0.005 wt % for Bi, 0.008 wt % for Se and Te, 0.01 wt % for Sb, 0.02 wt % for As and Sn, 0.03 wt % for Zr.

Oxygen (O) is involved in steel as an inevitable impurity. When excessive amounts of oxides are present in steel, ductility is reduced. In order to maintain ductility at very high levels, the content of O should be as low as possible. The upper limit of the amount of O should be 0.0020 wt % to avoid low-ductility issue, preferably 0.0015 wt %, most preferably 0.0015 wt %.

The steel described herein may be produced by an electric arc furnace. Tapping, deoxidation, alloying additions, and secondary metallurgy operations are typically carried out in a ladle furnace. The liquid steel may be continuously cast on a casting machine in round billets or cast in ingots that are processed to produce round bars of various diameters. Round bars may be re-heated (e.g., by a rotary hearth furnace) to a temperature between about 1200° C. and 1300° C. and hot pierced. The resulting hollows are hot rolled by a floating mandrel mill, or a multi-stand pipe mill with retained mandrel, or a Premium Quality Finishing (PQF®) mill, to yield tubular articles. After hot rolling, the tubular articles may be in-line heated, without cooling at room temperature, by an intermediate furnace for making temperature more uniform, and accurate sizing may be carried out by either a multi-stand stretch reducing mill or a sizing mill to bring the product to its required final size and tolerances.

Subsequently, the tubular articles may be cooled in air down to room temperature in a cooling bed. For a tubular articles having a final OD greater than about 16 inches, the tubular articles produced by the multi-stand pipe mill may be processed by a rotary expansion mill. For example, medium size (equal or less than 16 inches) tubular articles may be reheated by a walking beam furnace to a temperature between about 1150° C. and about 1250° C., expanded to the desired diameter by the expander mill at a temperature between about 1100° C. and about 1200° C., and in-line reheated before final sizing and cooling in still air. After cooling to room temperature, the tubular articles are then heat treated (normalization and tempering treatment) in a walking beam furnace or a continuous furnace. The furnace atmosphere may be either oxidizing (gas-fired furnace), chemically inert, or reducing in order to improve surface quality, thereby limiting or avoiding formation of an oxide scale.

Normalizing involves heating steel to a temperature where it is substantially entirely composed of austenite (e.g., heating to around at least 900° C.) and precipitates are dissolved, and then cooling in still air to ambient temperatures. The chemical composition range of the steel described herein is such that after normalizing, the steel consists essentially of martensite, bainite, or a combination thereof. Some amounts (less than 20% in volume) of ferrite and retained austenite can form at this stage.

In one example of a normalizing treatment, the tubular articles are first heated at about 20° C. per minute to between about 1040° C. and about 1080° C., to promote full transformation of ferrite into austenite, as well as dissolution of the precipitates in the austenitic phase.

The average cooling rate of a tubular article between 800° C. and 500° C. in still air during normalizing treatment depends is mostly affected by the wall thickness (WT): for WT between 2 mm and 15 mm, the cooling rate exceeds 1° C./s; for WT of about 20 mm, the cooling rate is around 0.8° C./s; for WT between 21 mm and 30 mm, the cooling rate is between about 0.4° C./s and 0.7° C./s. Thus, the tubular articles, after the normalizing treatment, have a final microstructure based on WT dimensions. For WT between 2 mm and 15 mm, the tubular articles have a microstructure of fully martensite. In one example, normalized tubular articles with WT between 2 mm and 15 mm have a martensite content in a range between about 90 vol % to about 100 vol % martensite, and may include ferrite, bainite, or both, neither of which cannot be detected by light microscopy. In another example, normalized tubular articles having a lean chemical composition and 15 mm<WT≦30 mm have a ferrite content of less than about 10 vol %, a bainite content of less than about 25 vol %, and a martensite content of at least 65 vol %. As used herein, “lean chemical composition” generally refers to a composition in which the content of the elements (those that are not explicitly indicated as residuals) is near the minimum value of the claimed range. Thus, normalized tubular articles with WT up to and including 30 mm have a predominantly martensite microstructure, with martensite exceeding 65 vol %. These tubular articles may also include bainite, ferrite, and less than 10 vol % of retained austenite.

In a final step of normalizing, the tubular articles are cooled in still air to room temperature, to allow transformation of austenite into martensite, bainite, or a combination thereof. Retained austenite and ferrite may be present in small amounts (less than about 20 vol %).

After normalizing, the tubular articles are tempered between about 700° C. and about 780° C. The tempering heat treatment promotes the transformation of martensite and bainite towards the stable form of ferrite and carbides, relieving internal stresses and thereby leading to more stable, predictable behavior of the steel in its most likely applications for high-temperature service. In addition to tempered martensite and tempered bainite, tempering promotes the formation and stabilization of a relatively uniform distribution of the small carbide and nitride particles desired for strengthening and achieves a desired state of precipitation. During tempering, retained austenite is decomposed into ferrite and carbides, and the final amount of retained austenite in the tempered tubular article is less than 2 vol %, less than 1 vol %, less than 0.5 vol %, or less than 0.1 vol %. That is, the tempered tubular article is substantially free of austenite, and may consist essentially of tempered martensite, tempered bainite, ferrite, carbides and nitrides, and any combination thereof.

The tubular articles may have an outer diameter between about 20 mm and about 624 mm and a wall thickness between about 2 mm and about 30 mm.

The tempered tubular articles are straightened to achieve the required straightness level.

Examples

Initial Testing. Laboratory-scale heats of alloy steels 1 to 6 with compositions shown in Table 1 were prepared in vacuum from pure charge materials. The tabulated compositions are the results of chemical analysis. Alloy steels 1 to 3 were formulated to evaluate the effect of adding Cu to a base 5Cr steel composition; alloy steels 4 to 6 were intended to evaluate the effect of varying C concentration with 2 wt % Cu added to the 5Cr steel base composition.

TABLE 1 Chemical compositions in wt % of main elements in alloy steels 1 to 6 * Alloy steel Cr Cu Mn Mo Si C Nb Ti 1 4.97 <0.01 0.40 0.50 0.22 0.130 <0.01 <0.01 2 4.97 1.00 0.38 0.50 0.21 0.055 <0.01 <0.01 3 4.93 1.97 0.37 0.50 0.20 0.082 <0.01 <0.01 4 4.57 1.89 0.27 0.50 0.26 0.086 0.05 0.04 5 4.66 1.90 0.28 0.50 0.26 0.045 0.05 0.04 6 4.65 1.94 0.28 0.50 0.25 0.120 0.05 0.05 * Additionally, all alloy steels contain S and P below 0.005 and 0.01 wt %, respectively, Ni content below 0.12 wt %.

The alloy steels were repeatedly arc-melted on a water-cooled Cu hearth plate to promote good mixing of constituent elements, and each was finally cast under vacuum into a Cu permanent mold. After hot-tops were removed, the size of the resulting ingots was 12.4 mm high, 25 mm wide, and 125 mm long. The ingots were homogenized in vacuum for 4 h at 1150° C. and hot rolled into 3-mm-thick plates. The hot-rolled plates were normalized for 1 h at 1050° C.

Specimens were cut from the hot-rolled plates and heat treated for 1 h at 300-800° C. to evaluate the tempering response of the alloy steels. The tempered specimens were mounted in epoxy and metallographically polished. Ten Vickers hardness (HV) indents were made on each polished specimen using an indenter load of 200 g. These results are tabulated in Table 2, and plotted in FIGS. 2A and 2B, in which the error bars represent one standard deviation. FIG. 2A shows variation of Vickers hardness with tempering temperature for alloy steels 1-3, with Cu additions of 0.1, 1.00, and 1.97 wt %, respectively. FIG. 2B shows variation of Vickers hardness with tempering temperature for alloy steels 4-6 containing about 2 wt % Cu and C levels of 0.086, 0.045, and 0.12 wt %, respectively.

TABLE 2 Variation of Vickers hardness measurements with tempering temperature Tempering Alloy Steel temperature Vickers hardness T (° C.) 1 2 3 4 5 6 300 414.7 376.2 410.4 413.5 373.0 423.6 350 426.7 383.4 403.5 427.0 377.7 416.5 400 424.5 386.0 415.5 416.5 384.9 415.7 450 423.4 393.8 428.3 437.4 405.3 432.3 500 429.0 401.9 470.1 460.8 433.4 465.3 550 442.7 404.4 441.8 435.6 418.3 451.3 600 344.0 291.7 320.0 341.2 321.5 347.7 650 266.6 243.4 285.3 284.2 292.0 290.9 700 231.6 212.0 237.8 259.4 264.4 265.5 750 191.6 186.8 202.1 221.0 207.4 222.2 800 214.9 287.3 384.0 347.9 290.3 393.4

Interpretation of the hardness results in FIG. 2A is complicated by the variations in C concentration, but each of alloy steels 1-3 experienced hardening due to tempering with their hardness peaks in the 500-600° C. range. The hardness of alloy steels 1 and 3 were nearly identical and higher than those of alloy steel 2 for tempering up to 750° C. Alloy steel 3 contained 1.97 wt % Cu, whereas alloy steel 1 had no intentional addition of Cu. However, the C concentration of alloy steel 1 was 0.13 wt % compared to that of 0.082 wt % in alloy steel 3. This difference in C concentration is relatively large and enough to influence the relative fractions of bainite or martensite and their hardness. A possible conclusion about the Cu addition is that it has a small effect on maintaining hardness at 650-750° C. The low C concentration of alloy steel 2, 0.055 wt %, may be at least partially responsible for its low hardness compared to alloy steels 1 and 3.

The results shown in FIG. 2B for the alloy steels 4-6 containing nominally 1.9 wt % Cu with the Nb+Ti additions show a similar trend. Alloy steels 4 and 6 with C concentrations of 0.086 and 0.12 wt %, respectively, had very similar hardness values for each tempering temperature up to 600-650° C. The hardness values of alloy steel 5, with a C concentration of 0.045 wt %, were the lowest of alloy steels 4-6 over this same temperature range. The hardness of alloy steels 4-6 were virtually identical for tempering at 650-750° C. However, over the same temperature range, the hardness of alloy steels 4-6 were higher than those of alloy steels 1-3. This suggests that the combination of Cu addition coupled with Nb+Ti additions promotes maintaining strength when tempering at 650-770° C.

FIGS. 2A and 2B also show that all of alloy steels 1-6 experienced higher hardness when tempered at 800° C. compared to 750° C. This may have been the result of the lower ferrite-austenite transformation temperature (A₁) being exceeded when tempering at 800° C. Exceeding that temperature could have resulted in fresh, untempered martensite forming in the microstructures, thereby increasing hardness relative to a fully tempered microstructure. The calculated A₁'s, shown in Table 3, are all in the 770-790° C. range, supporting this interpretation regarding increased hardness for tempering at 800° C.

TABLE 3 Calculated lower ferrite-austenite transformation temperatures (A₁) Alloy A₁ steel (° C.) 1 790 2 773 3 772 4 785 5 786 6 783

Optical micrographs of alloy steels 1-6 after normalizing at 1050° C. and tempering at 700° C. are shown in FIGS. 3A-3F, respectively. Microstructures of all of the alloy steels appear to be bainitic/martensitic. One clear difference between alloy steels 1-3 and alloy steels 4-6 is that the prior austenite grain sizes of the alloy steels without the Nb+Ti additions (alloy steels 1-3, shown in FIGS. 3A-3C, respectively, at a magnification of 100×) are considerably larger than those of the alloy steels containing these microalloying elements (alloy steels 4-6, shown in FIGS. 3D-3F, respectively, at a magnification of 500×), which are displayed at a 5 times higher magnification. The prior austenite grain size, about 200 μm, of alloy steel 1 (Cu<0.01) is the largest of all six alloy steels. The prior austenite grain sizes of the Cu-containing alloy steels 2 and 3 are similar to one another at about 150 μm, and slightly smaller than that of alloy steel 1. The prior austenite grain sizes of the alloy steels containing Nb+Ti (alloy steels 4-6) are all similar (10-20 μm), and much smaller than the grain sizes of the alloy steels without the Nb+Ti additions (alloy steels 1-3).

Tensile and creep test specimens were machined from the 3-mm-thick strips that were normalized and tempered at 700° C. The limited amount of metal available from the lab-scale heats required subsized specimens to be used for evaluations of tensile and creep properties. Six specimens were made of each alloy steel (e.g., specimens 1-1, 1-2, 1-3, 1-4, 1-5, and 1-6 were made of alloy steel 1).

Data from tensile tests run at room temperature, 550° C. and 600° C. are tabulated in Table 4 for three specimens of each alloy (e.g., 1-1, 1-2, 1-3 are specimens 1-3, respectively, of alloy steel 1). The tests were run using a cross-head speed of 1 mm/min and a clip-on extensometer to measure longitudinal strain in the gage section. Elevated temperature tests were equilibrated for 30 minutes prior to straining. The yield stress and peak stress of alloy steels 1-6 exceeded the minimum desired values of 450 MPa and 585 MPa, respectively. At room temperature, comparing the results from alloy steel 1 to those of alloy steel 3 indicates that the addition of 1.97 wt % Cu significantly increases yield and tensile stresses without compromising ductility. The combination of Cu addition with Nb+Ti did not significantly increase tensile stress beyond that of only the Cu addition. The data from the elevated temperature tests show a similar trend. However, these results can be affected by a low nitrogen content (below 0.005 wt % for all alloy steels) and a low C content (e.g. 0.045 wt % for alloy steel 5).

TABLE 4 Tensile test results Temper- Peak 0.2% Yield Alloy ature stress stress steel (° C.) (MPa) (MPa) 1-1 25 638 516 2-1 25 586 499 3-1 25 699 611 4-1 25 N.A. 607 5-1 25 672 609 6-1 25 735 618 1-2 550 398 373 2-2 550 379 359 3-2 550 437 417 4-2 550 419 398 5-2 550 437 427 6-2 550 449 426 1-3 600 339 320 2-3 600 332 317 3-3 600 375 359 4-3 600 359 335 5-3 600 383 367 6-3 600 383 363 N.A. - Broken, out of gage section

Data from stress-rupture testing at 550° C. of two additional specimens of alloy steels 1, 3, and 6 are summarized in Table 5A (e.g., 1-4 and 1-5 are specimens 4 and 5, respectively, of alloy steel 1). The tests were run without extensometers, so creep curves were not recorded. Comparing the results from specimens of alloy steel 1 to those of alloy steel 3 suggests the Cu addition is effective for increasing rupture life at 550° C. The results from alloy steel 6 suggest that rupture life is further enhanced by the Nb+Ti addition. The result for specimen 6-5 is particularly encouraging about the combined effects of adding both Cu and the microalloy elements Nb+Ti.

TABLE 5A Stress-rupture results (creep testing at 550° C.) Rupture Temper- Rupture Elonga- Alloy stress ature life tion steel (MPa) (° C.) (hours) (%) 1-4 215 550 78 24 1-5 205 550 167 22 3-4 215 550 172 19 3-5 205 550 278 25 6-4 215 550 202 22 6-5 205 550 1,084 36

Hot Ductility Testing.

Specimens of normalized sheets of alloy steels 1 and 3 were tested to assess how the 1.97 wt % Cu addition affected high temperature workability. The specimens used for this test had dimensions of 2.5 mm thickness×5.6 mm width×75 mm length. They were gripped in a Gleeble machine using a free span of 25 mm and a fixed axial load of 24 kg. Specimen temperatures were monitored and controlled by thermocouples attached at the mid-point of the free-span lengths as is usual for this type of Gleeble test. The specimens were programmed to heat to 1200° C. at a rate of 100° C./s and hold at 1200° C. for 10 s. This was immediately followed by heating at 1° C./s to 1300° C., holding at 1300° C. for 10 s, and then heating at 1° C./s to 1350° C. The results from heating specimens with this procedure in air are shown in FIGS. 4A and 4B.

FIG. 4A shows variations of specimen temperature with time during Gleeble testing. FIG. 4B shows the variations of specimen temperature from 1200° C. and up. The actual temperature of both specimens followed the programmed temperature well. The specimen of alloy steel 1 reached a maximum temperature of 1317° C. before it failed due to the 24-kg axial load. The specimen of alloy steel 3 reached a maximum temperature of 1314° C. This result indicates that the addition of 2 wt % Cu to the 5Cr base alloy did not have a significant effect on its nil-ductility temperature (NDT).

FIG. 4B shows variations of axial displacement with temperature during Gleeble testing. The variations of axial displacements of the two specimens due to the 24-kg axial loading are shown in FIG. 4B for the same portions of the heating programs. A small vertical offset was used to plot the curves in FIG. 4B so that both could be clearly viewed. The displacement response of both heats was virtually identical. This suggests that the very-high-temperature ductility of alloy steels 1 and 3 were similar.

The specimens of alloy steels 1 and 3 were both heavily oxidized due to exposure to air during the heating. Necking developed in and appeared to be similar for both specimens. There is no clear indication of incipient melting or unusual liquation in alloy steel 3.

Transformation Behavior During Continuous Cooling.

The A_(C1) temperatures of specimens of alloy steels 1-3 and 6 were estimated using a Gleeble machine following the critical steps of the procedure described in ASTM A1033-10. The tests were run under vacuum. Before any measurements, the specimens were given stabilization treatments in the Gleeble of heating to 650° C. at 10° C./s, holding at 650° C. for 10 min, and then cooling to room temperature at a maximum rate of 20° C./s. The specimens used for the Gleeble tests were 2.5-mm thick 5.5-mm wide and about 75-mm long.

The A_(C1) temperatures were determined by heating specimens to 750° C. at 10° C./s, then from 750-850° C. at 28° C./min, and finally from 850-900° C. at 10° C./s. This procedure was repeated three times with each specimen using cooling rates of 0.1° C./s, 1° C./s, and 10° C./s. The A_(C1) temperatures and the transformation temperatures on cooling were estimated graphically as outlined in ASTM A1033-10.

The A_(C1) temperatures are given in Table 5B. For each alloy there was a tendency for the A_(C1) of the first heating cycle to be higher than those of the following heating cycles. Comparing the measurements from alloy steels 1 and 2 indicates that the addition of 1 wt % Cu reduced the A_(C1) by about 10-20° C. Further addition up to 2 wt % Cu in alloy steel 3 produced no significant lowering of A_(C1) beyond that of alloy steel 2. The carbon concentrations of alloy steels 1-3 varied from 0.055-0.13 wt %, which would have a small effect on the A_(C1). The carbon concentrations of alloy steels 1 and 6 were nearly identical: 0.13 compared to 0.12 wt % C. The A_(C1) of alloy steel 6 is 790° C. indicating that the 2 wt % Cu decreased this transformation temperature by about 10° C.

TABLE 5B Measured A_(C1) temperatures Alloy Ac₁ steel (° C.) 1-1 806 1-2 802 1-3 799 2-1 791 2-2 781 2-3 780 3-1 785 3-2 782 3-3 782 6-1 792 6-2 787 6-3 791

The transformation temperatures on cooling are tabulated in Table 6. One effect Cu appears to have on the transformation behavior is the suppression of ferrite formation at 1° C./s cooling and 2 wt % Cu.

TABLE 6 Measured transformation temperatures Transformation Temperatures (° C.) at indicated cooling rates 10° C./s 1° C./s 0.1° C./s Transformation Alloy Bainite Ferrite Bainite Ferrite Bainite steel start start start start start 1 473 781 478 802 443 2 492 775 498 790 461 3 455 Not observed 472 739 458 6 444 Not observed 472 754 405

Summary of Initial Testing.

Alloy steels 1-6 were tensile tested at room temperature, 550° C., and 600° C. All the room temperature results exceeded 450 MPa yield strength and 585 MPa tensile strength. The addition of 2 wt % Cu increased room temperature strength about 20% over alloy steel 1. Although the amounts were lower, a strength advantage was maintained at the higher temperatures. Some steel compositions do not show the ferrite transformation start for cooling rates of 1° C./s. Slower cooling rates are required for the ferrite microconstituent to form.

Stress-rupture testing at 550° C. indicated that the combined additions of Cu+Nb+Ti more than doubled rupture life compared to alloy steel 1.

Nil-ductility temperature testing on a Gleeble machine suggested that Cu additions up to 2 wt % do not produce liquation or incipient melting that degrades workability.

Lower critical ferrite-to-austenite temperatures, A_(C1) temperatures, were measured using a Gleeble machine following the procedures of ASTM 1033-10. The results showed that, compared to alloy steel 1, a Cu addition of 1 wt % reduced the A_(C1) temperature by 10-20° C. Increasing the Cu addition up to 2 wt % had no significant further effect on reducing A_(C1) temperature.

Austenite transformation temperatures during cooling from 900° C. were also measured on a Gleeble machine. For cooling at 10° C./s and at 0.1° C./s, the Cu additions had no significant effect on transformation behavior. For cooling at 1° C./s, both alloy steel 1 and 2 deviated from the expected behavior by indications of ferrite formation.

Additional Testing.

An 80 kg experimental heat of alloy steel 7 was prepared. The composition of alloy steel 7 is shown in Table 7, along with the composition of the closest matching alloy from the initial testing (alloy steel 4).

TABLE 7 Chemical compositions in wt % of main elements of alloy steel 7 compared to alloy steel 4 Alloy steel C Mn Si Mo Cr Ni Cu S Nb Ti N 7 0.071 0.38 0.26 0.53 4.62 0.12 1.71 0.0029 0.054 0.051 0.0097 4 0.086 0.27 0.26 0.50 4.57 <0.12 1.89 <0.0050 0.050 0.040 <0.0050

The cast ingot was hot rolled down to a thickness of 20 mm (finish rolling temperature of 1010° C.), and then cut into 8 pieces. One plate was left as-rolled (asR) and the remaining ones were normalized at 1050° C.; one of these was left only normalized. Another normalized plate was tempered at 700° C. for 1 h; three other plates were tempered at 750° C. for 1 h. Finally, the two remaining slabs were subject to a second normalization, at a temperature of 950° C., then tempered for 1 h at 750° C. Table 8 summarizes the samples that have been produced from the experimental heat.

TABLE 8 Alloy steel 7 heat treated plates Size Heat Treatment Details Alloy (mm × mm × Normalized Normalized Tempered Tempered steel mm) Heat treatment at 1050° C. at 950° C. at 700° C. at 750° C. 7-A-1 155 × 260 × Normalizing X X 18.5 and Tempering 7-A-2 155 × 273 × Normalizing X X 18.5 and Tempering 7-A-3 155 × 273 × Normalizing X X 18.5 and Tempering 7-A-4 155 × 273 × Normalizing X X 18.5 and Tempering 7-B-1 155 × 273 × second X X X 18.5 normalization and Tempering 7-B-2 155 × 273 × second X X X 18.5 normalization and Tempering 7-B-3 155 × 273 × Normalizing X 18.5 7-B-4 155 × 273 × asR (no heat Dilatometry 18.5 treatment);

Continuous cooling transformations were measured by dilatometry, starting from the as-rolled material. Austenitization was performed at 920° C. (75° C. above A_(C3)) for 12 min. Prior austenitic grain size was measured by the lineal intercept method (ASTM E112) after Winsteard etching, resulting in 9.3 μm mean lineal intercept (MLI). The continuous cooling transformation (CCT) diagram of alloy steel 7 is shown together with a comparison with alloy steel 4 (which has a similar composition, except the addition of Ti and Nb) in FIG. 5A. Experimentally measured transformation points (1-99% transformation) for alloy steel 7 are indicated as empty round dots 510, 511, 512, 513, and 514; solid lines delimit transformation ranges. Solid dots and dashed lines indicate transformation temperatures and ranges for alloy steel 4. HV hardness for alloy steel 7 is indicated at the bottom of each cooling curve inside a circle. “M,” “F,” and “P” in FIG. 4A refer to martensite, ferrite, and pearlite, respectively.

As shown in FIG. 5A, alloy steel 7 exhibits a fully martensitic structure for cooling rates greater than 1° C./s. No significant difference in terms of hardness and microstructure examined by light microscopy was seen for cooling rates greater than or equal to 1° C./s. FIG. 5B shows an optical micrograph of a specimen of alloy steel 7 (381 HV₁₀) cooled at a rate of 1° C./s. A fully martensitic structure (partially self-tempered) was observed. FIG. 5C shows an optical micrograph of a specimen of alloy steel 7 (375 HV₁₀) cooled at a rate of 60° C./s. FIG. 5D shows an optical micrograph of a specimen of alloy steel 7 after cooling at 0.5° C./s. For a cooling rate of 0.5° C./s, a small amount of ferrite was seen (6% in volume), whereas the remaining microstructure still exhibits hardness and morphology typically of martensite (indicated by “M”) formed at cooling rates of about 1° C./s. Small amounts of bainite may be formed together with ferrite (indicated by “F”) at a cooling rate of less than about 0.5° C./s (e.g., for tubular articles with a wall thickness of about 30 mm) for the leaner chemical composition (e.g., compositions low in C, Cr and Mo). However, at cooling rate of about 0.5° C./s, even in the case of a lean composition, the amount of tempered bainite present, based on CCT diagrams and metallurgical models, is less than 25%.

Phase amount after normalization of alloy steel 7 was measured as a function of cooling rate by point counting (according to ASTM E562) after metallographic etching with Nital reagent (2 vol % concentrated nitric acid in ethanol). Results are shown in FIG. 7.

Retained Austenite after Normalization and Tempering.

Austenite was not detected by X-ray diffraction technique. According to the sensitivity of the applied technique the maximum amount of retained austenite is below 2 vol %.

Mechanical Properties after Normalization and Tempering.

Room temperature mechanical properties (yield stress, YS, and ultimate tensile stress, UTS) on differently treated normalized and tempered samples were measured by longitudinal tensile test and impact Charpy-V notch (CVN) test in both longitudinal (L) and transverse (Tr) direction. Results are summarized in Table 9, along with those of alloy steels 3 and 6. Alloy steel 3 has a composition similar to alloy steel 7, except the presence of Nb and Ti as microalloying elements. Alloy steel 6 includes Nb and Ti.

TABLE 9 Strength properties of alloy steel 7 compared to alloy steels 3 and 6. Alloy Normalizing Tempering YS_(L) UTS_(L) El._(L) CVN_(L) CVN_(Tr) steel temperature (° C.) (° C.; 1 hour) (MPa) (MPa) (%) (J) (J) HV 7 1050° C. 700° C.; 1 hour 646 725 24 260 212 243 7 1050° C. 750° C.; 1 hour 563 658 26 293 266 218 7 1050° C. + 700° C.; 1 hour 526 639 27 311 245 207 950° C. 3 1050° C. 700° C.; 1 hour 611 699 18 N.A. N.A. 238 6 1050° C. 700° C.; 1 hour 618 735 14 N.A. N.A. 265 N.A.—Not available (CVN specimens could not be prepared from the available thin material)

Precipitates Characterization.

Precipitate phases at equilibrium in alloy steel 7 were calculated using the JMatPro software package. Calculations show that precipitation of M_((1,2))C and M₂₃C₆ carbides, as well as metallic ε-Cu is to be expected. Computer simulations indicate that MC type carbides are composed almost entirely of Ti, M₂C of Nb and M₂₃C₆ of Cr, Mo and Mn. Phase composition at equilibrium as a function of temperature is shown FIG. 6.

Precipitates were studied by transmission electron microscopy on the normalized (1050° C.) and tempered (750° C.) alloy steel 7 material.

FIG. 8A shows a TEM image of a thin film cut from the bulk plate. Copper precipitates 110 appear as dark, somewhat round spots within the metallic matrix grains. These metallic ε-Cu precipitates were identified by energy-dispersive X-ray spectroscopy (EDS). FIG. 8B shows an EDS spectrum of the matrix, and FIG. 8C shows an EDS spectrum of a copper precipitate. A sharp increase in the Cu-Kα peak intensity is seen in FIG. 8C (e.g., when the electron beam is focused on a precipitate particle). Cu incoherent precipitates have been measured and appear to be 31±7 nm in size. Assuming concentration to be equal to the equilibrium value at tempering temperature (750° C.), this value corresponds to 1.3×10¹⁵ cm⁻³.

Carbide precipitation was characterized by TEM imaging on an extraction replica. Carbides 112 were found to appear along grain boundaries and within laths/grains, as shown in the TEM image of FIG. 9A. Spots on which energy-dispersive spectroscopy (EDS) was performed are labeled by type, with type i (Cr, Fr, Mo), ii (Cr, Fe, Mo, Si), and iii (Cr, Fe, Mo, Nb, Ti) spots identified by EDS as corresponding to M₃C, M₂₃C₆, and M_((1,2))(C,N) carbo-nitride precipitation.

An example of an EDS spectrum from a M₂₃C₆ type carbide is shown in FIG. 9B. Carbide size distribution measured by digital image analysis over non-overlapping particles yielded a mean size of 59±5 nm. FIG. 9B shows a surface area diffraction pattern (SAD pattern) that is obtained using transmission electron microscopy. Electrons hitting a crystal, such as a precipitate that is encased in the extraction replica, are diffracted by the crystal lattice at specific angles that are correlated to distance between atomic planes in the crystal. On the right, the intensity of the diffracted signal as a function of the scattering vector, a spatial coordinate that is function of the angle between the incoming beam and outgoing beam from the specimens, can be found. FIG. 9B shows the measured average intensity. The thicker lines and the dotted lines represent the positions at which peaks from different crystal structures (in this case, two different carbide species) are expected to be found. Correspondence between the thicker lines or dotted lines and the average intensity pattern, combined with chemical analysis performed by means of EDS (energy dispersive X-ray spectroscopy) helps identifying the type of precipitate.

Chromium Content in Solid Solution.

After tempering, alloy steel 7 is constituted of an alpha ferrite matrix and precipitates (mainly carbides). FIG. 11 depicts the Cr content in wt %, calculated under equilibrium conditions, in solid solution in the ferritic matrix after tempering at various temperatures. No more than 4 wt % of Cr is stable in solid solution at room temperature; the rest of Cr goes into precipitates.

Dislocation Density.

Dislocation density was measured by X-ray diffraction (XRD) line profile analysis on alloy steel 7 samples after normalization and tempering at two different temperatures (700° C. and 750° C.). XRD patterns are shown in FIGS. 10A and 10B. Arrows indicate residual austenite peaks, whose total amount is on the order of 2 wt %. In some embodiments, the total amount of residual austenite is less than 1 wt %.

Mean crystalline domain size was estimated to be 290 and greater than 400 nm (beyond the technique's resolution), respectively, for the samples tempered at 700° C. and 750° C. Dislocation density was measured to be low in both samples, being 4.6×10¹⁴ m⁻² and 4.2×10¹⁴ m⁻², respectively.

Microstructural Contributions to Mechanical Properties.

Yield strength was found to vary from 646 to 563 MPa when tempering was performed at 700° C. and 750° C., respectively. Dislocation density is known to contribute to the final yield strength according to the following formula: σ_(γ)≈0.5·G·b·ρ^(0.5), where G is the shear modulus of the material, b is the Burgers vector length, and ρ is the dislocation density (T. Gladman, “Physical Metallurgay of Microalloyed Steels,” Maney Pub. 1997). According to this model, the estimated Δσ_(γ) due to dislocations is estimated to be around 10 MPa. Precipitates contribute to the final yield strength according to the following formula: σ_(γ)=(10.8·fv^(0.5)/<D_(p)>)ln(<D_(p)>/6.125·×10⁻⁴), where <D_(p)> is the average precipitate size and fv the volume fraction (Gladman). The estimated Δσ_(γ) due to Cu precipitation is 56 MPa. This value is in reasonable agreement with the 94 MPa increase observed in the initial testing upon 2% Cu addition (considering that small composition differences also play a role). Carbide precipitation yields another, theoretical 82 MPa increase.

As precipitate characterization by means of TEM was only performed on the sample tempered at 750° C., the latter two results were estimated in the following fashion: precipitate size was assumed to keep constant, whereas volume fraction was assumed to be equal to the equilibrium value at tempering temperature.

The sum of all effects results in a yield strength difference of about 150 MPa, larger than the measured value. This can be explained by the fact that precipitates are likely to be present in smaller size when tempering is performed at lower temperature, because of their slower growth rate. For example, the correct yield strength is obtained by assuming <D_(Cu)>=20 nm (down from 31) and <D_(carbides)>=37 nm (down from 59).

Summary of Additional Testing.

An 80 kg heat of 5Cr-2Cu—Nb—Ti steel (alloy steel 7) was cast and hot rolled to a 20-mm-thick plate by pilot mill (finish rolling temperature of 1010° C.). The chemical composition of alloy steel 7 (0.07C-0.38Mn-0.26Si-0.53Mo-4.6Cr-0.12Ni-1.7Cu-0.05Nb-0.05Ti-0.01N) was obtained on the basis of the results on creep properties of alloy steels 1-6.

A critical cooling rate of 0.7° C./s was measured by dilatometry for a fully martensitic structure. Microstructural features such as precipitates (copper incoherent precipitates and carbides) were measured by TEM; and dislocation density and crystalline domain size by XRD line profile analysis. Room temperature mechanical properties were related to these microstructural features by physical metallurgy models. Cu was estimated to contribute for about 60 MPa increase in yield strength, whereas fine carbides were estimated to contribute another 80 MPa.

A number of embodiments of the invention have been described. Nevertheless, it will be understood that various modifications may be made without departing from the spirit and scope of the invention. Accordingly, other embodiments are within the scope of the following claims. 

What is claimed is:
 1. A tubular article formed from high temperature steam oxidation resistant and high temperature creep resistant alloy steel, the chemical composition of which comprises, by weight: C of 0.06 to 0.15 wt %; Si of 0.1 to 0.5 wt %; Mn of 0.2 to 0.6 wt %; Ni of 0.05 to 0.4 wt %; Cr of 4.5 to 6.0 wt %, Cu of 1.0 to 2.0 wt %; Ti of 0.04 to 0.08 wt %; Nb of 0.01 to 0.06 wt %; Mo of 0.45 to 1.2 wt %; N of 0.008 to 0.05 wt %; Al of less than or equal to 0.01 wt %; Zr of less than or equal to 0.01 wt %; Co of less than or equal to 3.0 wt %; V of less than or equal to 0.07 wt %; W of less than or equal to 3.0 wt %; P of less than or equal to 0.015 wt %; S of less than or equal to 0.003 wt %; and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or equal to 0.1 wt %; and a balance comprising Fe and impurities; and wherein the steel comprises: fine carbides, nitrides, or both; and copper precipitates, and wherein the steel has a final microstructure comprising tempered martensite, tempered bainite, or a combination thereof; and wherein the steel contains less than 2 vol % residual austenite.
 2. The tubular article of claim 1, wherein the steel contains less than 1 vol % residual austenite.
 3. The tubular article of claim 1, wherein less than 4 wt % of the total Cr amount is stable in solid solution at room temperature.
 4. The tubular article of claim 1, wherein the steel is comprising maximum 25 percent bainite.
 5. The tubular article of claim 1, wherein wt % Ta+wt % Nb<0.1 wt %.
 6. The tubular article of claim 1, wherein the copper precipitates have a grain size of less than 40 nm.
 7. The tubular article of claim 1, wherein the fine carbides have a mean diameter of less than 200 nm.
 8. The tubular article of claim 7, wherein the fine carbides have a size distribution with a mean size of 65 nm or less.
 9. The tubular article of claim 1, wherein the steel has a yield strength greater than 450 MPa.
 10. The tubular article of claim 1, wherein the steel has a yield strength between 570 MPa and 640 MPa.
 11. The tubular article of claim 1, wherein the steel has an ultimate tensile strength greater than 585 MPa at room temperature.
 12. The tubular article of claim 1, wherein the steel has a hardness of 200 to 265 Vickers.
 13. The tubular article of claim 1, wherein the steel has a creep strength of about 70 MPa for 100,000 h rupture at 600° C.
 14. The tubular article of claim 1, wherein the steel has a minimum absorbed energy in Charpy V-notch impact test of 150 J at 0° C.
 15. The tubular article of claim 1, wherein the steel consists essentially of: C of 0.06 to 0.15 wt %; Si of 0 to 0.5 wt %; Mn of 0.2 to 0.6 wt %; Ni of 0 to 0.4 wt %; Cr of 4.0 to 6.0 wt %: Cu of 1.0 to 2.0 wt %; Ti of 0.04 to 0.08 wt %; Nb of 0.01 to 0.06 wt %; Mo of 0.5 to 1.2 wt %; N of 0.008 to 0.02 wt %; Al of less than or equal to 0.01 wt %; Zr of less than or equal to 0.01 wt %; Co of less than or equal to 2.0 wt %; V of less than or equal to 0.07 wt %; W of less than or equal to 2.0 wt %; P of less than or equal to 0.015 wt %; S of less than or equal to 0.003 wt %; and Ca, Ta, Mg, Se, Te, B and Bi in contents of less than or equal to 0.1 wt %; and a balance comprising Fe and impurities.
 16. The tubular article of claim 1, wherein the alloy steel consists essentially of: C of 0.06 to 0.15 wt %; Si of 0 to 0.5 wt %; Mn of 0.2 to 0.6 wt %; Ni of 0 to 0.4 wt %; Cr of 4.0 to 6.0 wt %; Cu of 1.0 to 2.0 wt %; Ti of 0.04 to 0.08 wt %; Nb of 0.01 to 0.06 wt %; Mo of 0.5 to 1.2 wt %; N of 0.008 to 0.02 wt %; Al of less than or equal to 0.01 wt %; Zr of less than or equal to 0.01 wt %; Co of less than or equal to 2.0 wt %; V of less than or equal to 0.07 wt %; W of less than or equal to 2.0 wt %; P of less than or equal to 0.015 wt %; Ta of less than or equal to 0.1 wt %; S of less than or equal to 0.005 wt %; Se of less than or equal to 0.005 wt %; and a balance comprising Fe and impurities.
 17. The tubular article of claim 1, wherein the steel has a final microstructure consisting essentially of martensite, bainite, or a combination thereof.
 18. The tubular article of claim 1, wherein the fine carbides comprise fine titanium carbides.
 19. The tubular article of claim 1, wherein the fine carbides comprise fine niobium carbides.
 20. The tubular article of claim 1, wherein the tubular article has a wall thickness between 2 mm and 30 mm.
 21. The tubular article of claim 1, wherein the outer diameter of between 20 mm and 624 mm.
 22. A method of forming a tubular article, the method comprising: providing an ingot or continuous casting of an alloy steel by combining the following added elements: C of 0.06 to 0.15 wt %; Si of 0 to 0.5 wt %; Mn of 0.2 to 0.6 wt %; Ni of 0 to 0.4 wt %; Cr of 4.0 to 6.0 wt %, Cu of 1.0 to 2.0 wt %; Ti of 0.04 to 0.08 wt %; Nb of 0.01 to 0.06 wt %; Mo of 0.5 to 1.2 wt %; N of 0.008 to 0.02 wt %; Al of less than or equal to 0.01 wt %; Zr of less than or equal to 0.01 wt %; Co of less than or equal to 2 wt %; V of less than or equal to 0.07 wt %; W of less than or equal to 2 wt %; P of less than or equal to 0.015 wt %; Ta of less than or equal to 0.1 wt %; S of less than or equal to 0.005 wt %; Se of less than or equal to 0.005 wt %; Ca of less than or equal to 0.1 wt %; Mg of less than or equal to 0.1 wt %; Se of less than or equal to 0.1 wt %; Te of less than or equal to 0.1 wt %; B of less than or equal to 0.1 wt %; Bi of less than or equal to 0.1 wt %; and a balance comprising Fe and impurities; forming the alloy steel into a raw tubular article; sizing the raw tubular article to yield a sized raw tubular article; cooling the sized raw tubular article to ambient temperature; normalizing the sized raw tubular article to yield a normalized tubular article having a microstructure comprising: martensite, bainite, or a combination thereof; and less than 20 vol % austenite; and tempering the normalized tubular article to yield a tempered tubular article comprising: alpha ferrite matrix, fine carbides and nitrides; and less than 2 vol % austenite.
 23. The method of claim 22, wherein normalizing the sized raw tubular article comprises: heating the sized raw tubular article to yield a heated sized raw tubular article having a microstructure consisting essentially of austenite; and cooling the heated sized raw tubular article to ambient temperature to yield the normalized tubular article.
 24. The method of claim 23, wherein heating the sized raw tubular article comprises heating the sized raw tubular article to a temperature in a range between 1040° C. and 1080° C. for at least 20 minutes.
 25. The method of claim 24, further comprising heating the sized raw tubular article at a rate between 15° C./min and 25° C./min.
 26. The method of claim 23, wherein cooling the heated sized raw tubular article comprises cooling in still air.
 27. The method of claim 23, wherein the wall thickness of the sized raw tubular article is between 2 mm and 15 mm, and the average cooling rate of the heated sized raw tubular article between 800° C. and 500° C. exceeds 1° C./s.
 28. The method of claim 23, wherein the wall thickness of the sized raw tubular article is about 20 mm, and the average cooling rate of the heated sized raw tubular article between 800° C. and 500° C. is about 0.8° C./s.
 29. The method of claim 23, wherein the wall thickness of the sized raw tubular article is between 21 and 30 mm, and the average cooling rate of the heated sized raw tubular article between 800° C. and 500° C. is between 0.4° C./s and 0.7° C./s.
 30. The method of claim 22, wherein tempering the normalized tubular article comprises heating the normalized tubular article to a temperature between 700° C. and 780° C. for at least 20 minutes, thereby decomposing the retained austenite to yield ferrite and carbides.
 31. The method of claim 30, further comprising straightening the treated tubular article.
 32. The tubular article of claim 1, wherein the chemical composition is free of at least one of Al, Zr, Co, V, W, Ca, Ta, Mg, Se, Te, B, or Bi. 